Flat Steel Product and Method for the Production Thereof

ABSTRACT

The invention provides a reliably manufacturable flat steel product based on a Fe3Al alloy and a method that allows the production of such flat steel products. For this purpose, the flat steel product is made of a steel which comprises (in % by weight) Al: 12-20%, Ti: 0.2-2%, B: 0.1-0.6%, in each case at least one element from the group “Cr, C, Mn, Si, Nb, Ta, W, Zr, V, Mo, Ni, Cu, Ca, rare earth metals, and Co” unavoidable impurities. The method according to the invention specifies that a molten steel with the stated composition is cast as a precursor in the form of a slab, thin slab or a cast strip, the precursor is then hot rolled into a hot strip.

The invention relates to a Fe—Al—Ti—B-based flat steel product and a method for producing such a flat steel product.

Whenever information is given in the present text on levels of certain elements in an alloy, if nothing else is expressly stated, these levels always refer to the weight (“% by weight”) or the mass (“% by mass”) of the alloy in question. By contrast, unless otherwise explicitly stated, information on microstructural components always refers to the volume occupied by the respective microstructure (“% by volume”).

When the term “flat steel products” is used in the present text, it refers to rolled products formed as strips, sheet metal or the cuts and blanks derived therefrom. The flat steel products according to the invention are in particular heavy plates with a typical sheet thickness of 6-200 mm, or hot-rolled strips or strips with a typical sheet thickness of 1.5-6 mm.

Steels of the type in question are characterised by TiB₂ precipitates embedded in a Fe₃Al matrix. As a result of this characteristic, such steels are low in density and, consequently, low in weight. These properties, which are interesting for many applications, are balanced in known materials of the type in question here by high brittleness up to high temperatures and insufficient strength at temperatures above 500° C.

The fundamental potential of materials based on the intermetallic phases Fe₃Al and FeAl was recognized around 100 years ago. Since then, there have been repeated attempts to develop materials based on the Fe₃Al phase in particular. However, it has not yet been possible to produce strip and sheet products from these materials.

A typical example of such attempts is described in EP 0 695 811 A1. The heat-resistant iron-based alloy presented therein is composed of the general formula Fe_(x)Al_(y)C_(z), where (in at. %) variable y is 1%≤y≤28% and variable z is ≤24%, and variable x is determined based on a diagram depending on the respective C and Al content of the steel. It is mentioned as an aside that the steel may contain more than forty other elements, including TiB₂, whereby each of these elements has a content allowance of 0.1-2% by atomic weight. How steel designed in such a way can be processed into flat steel products is not specified.

Other research has focused on producing Fe₃Al casting alloys based on boride-strengthened alloys. The results of this work can be found in the articles “Microstructure and mechanical properties of Fe₃Al-based alloys with strengthening boride precipitates” by Krein, R., et al. in Intermetallics, 2007. 15 (9): p. 1172-1182, and “The influence of Cr and B additions on the mechanical properties and oxidation behaviour of L2₁-ordered Fe—Al—Ti-based alloys at high temperatures” by Krein, R. and M. Palm in Acta mater., 2008.56 (10): p. 2400-2405., “L2₁-ordered Fe—Al—Ti alloys” by Krein, R., et al. in Intermetallics, 2010. 18: p. 1360-1364. Accordingly, based on the system, it is possible to produce Fe—Al—Ti—B fine-grained alloys whose microstructure consists of a Fe₃Al matrix with very small borides (<1 μm) along the grain boundaries. The compositions of the alloys are chosen so that the Fe₃Al phase primarily precipitates, while the borides are precipitated in the (residual) eutectic. In this way, the borides increase strength, improve ductility and fix the grain size of the Fe₃Al matrix.

For example, it is known from Li, X., P. Prokopcakova and M. Palm “Microstructure and mechanical properties of Fe—Al—Ti—B alloys with addition of Mo and W”, Mat. Sci. Eng., 2014, A 611: p. 234-241 that Fe—Al—Ti—B casting alloys can also be modified by adding further elements. In particular elements that increase the D0₃/B₂ transition temperature are considered here. In addition, Mo promotes the formation of complex borides, so that no more TiB₂ is formed.

Given the prior art described above, the object was to provide a flat steel product based on a Fe₃Al alloy and a method that allows reliable production of such flat steel products.

Regarding the flat steel product, the invention has achieved this object with a flat steel product created according to claim 1.

For reliable production of such flat steel products, the invention proposes the method specified in claim 11.

Advantageous embodiments of the invention are defined in the dependent claims and, like the general concept of the invention, are explained in detail in the following.

A flat steel product according to the invention is characterised in that it is made of a steel which consists of (in % by weight):

-   -   Al: 12-20%     -   Ti: 0.2-2%     -   B: 0.1-0.6%     -   , as well as optionally one or more elements from the group “Cr,         C, Mn, Si, Nb, Ta, W, Zr, V, Mo, Ni, Cu, Ca, rare earth metals,         Co” at the following levels:         -   N: up to 0.1%         -   Cr: up to 7%         -   C: up to 0.15%         -   Mn: up to 2%         -   Si: 0.05-5%         -   Nb, Ta, W: up to 0.2% in total         -   Zr: up to 1%         -   V: up to 1%         -   Mo: up to 1%         -   Ni: up to 2%         -   Cu: up to 3%         -   Ca: up to 0.015%         -   Rare earth metals: up to 0.2%         -   Co: up to 1%,     -   residual iron and unavoidable impurities, wherein S levels of up         to 0.03% by weight and P levels of up to 0.1% by weight can be         attributable to the unavoidable impurities.

It is critical to the invention that the ratio % Ti/% B formed from the Ti content % Ti and the B content % B of the steel is

0.33≤% Ti/% B≤3.75

and the structure of the steel or the resulting flat steel product consists of 0.3-5% by volume TiB₂ precipitates that are embedded in a structural matrix with at least 80% by volume Fe₃Al.

Due to its special composition, the boride-strengthened Fe₃Al alloy that a flat steel product according to the invention is made of exhibits strength above 500° C. and ductility that is markedly improved over conventional alloys of this type known from the prior art. At the same time, the parameters of the production of the flat steel product according to the invention from such a composite steel are set according to the invention in such a way that microstructure optimisation is achieved, by which the properties of a flat steel product according to the invention are further optimised.

For this, the method according to the invention for producing a flat steel product according to the invention comprises the following steps:

-   -   a) Melting a steel that consists of (in % by weight):         -   Al: 12-20%         -   Ti: 0.2-2%         -   B: 0.1-0.6%         -   , as well as optionally one or more elements from the group             “Cr, C, Mn, Si, Nb, Ta, W, Zr, V, Mo, Ni, Cu, Ca, rare earth             metals, Co” at the following levels:             -   N: up to 0.1%             -   Cr: up to 7%             -   C: up to 0.15%             -   Mn: up to 2%             -   Si: 0.05-5%             -   Nb, Ta, W: up to 0.2% in total             -   Zr: up to 1%             -   V: up to 1%             -   Mo: up to 1%             -   Ni: up to 2%             -   Cu: up to 3%             -   Ca: up to 0.015%             -   Rare earth metals: up to 0.2%             -   Co: up to 1%         -   residual iron and unavoidable impurities, wherein S levels             of up to 0.03% by weight and P levels of up to 0.1% by             weight can be attributed to the unavoidable impurities, and         -   wherein the ratio % Ti/% B formed from the Ti levels % Ti             and the B levels % B of the steel is

0.33≤% Ti/% B≤3.75;

-   -   b) Casting the steel melt into a precursorin the form of a slab,         thin slab or cast strip;     -   c) Hot rolling the precursor into a hot-rolled hot strip,         wherein the precursor has a hot rolling start temperature of         1000-1300° C. at the start of the hot rolling process, and the         hot rolling end temperature is at least 850° C.;     -   d) Coiling the hot strip at a coiling temperature between room         temperature and 750° C.

Aluminium is contained in the steel flat product according to the invention in levels of 12-20% by weight. At Al levels of at least 12% by weight, in particular more than 12% by weight, the intermetallic iron aluminide phase Fe₃Al forms, which is the main component of the microstructure of a flat steel product according to the invention. The high Al levels lead to reduced density, accompanied by reduced weight, high resistance to corrosion and oxidation, as well as high tensile strength. However, excessive Al levels would make it difficult to cold form the steels according to the invention. Excessive Al levels would also result in deteriorated weldability by forming a stable weld slag in the welding process, and increased electrical resistance in resistance welding. For these reasons, the Al level of steel according to the invention is limited to a maximum of 20% by weight, in particularof 16% by weight.

Ti and B form titanium borides in the steel according to the invention, which produce a fine microstructure, an increased yield strength, higher ductility, a higher modulus of elasticity and increased resistance to wear. In order for these effects to be achieved, Ti levels of at least 0.2% by weight, in particular at least 0.4% by weight, and B levels of at least 0.10% by weight, in particular at least 0.15% by weight are required.

It is essential for the invention that the Ti levels % Ti are matched to the B levels % B of the steel, so that the ratio % Ti/% B, i.e. the quotient of the Ti content % Ti as a dividend and the B content % B as a divisor, is 0.33 to 3.75, in particular 0.5-3.75 or 1.0-3.75. The risk of FeB formation is reduced when the ratio % Ti/% B is at least 0.33. Otherwise, the low melting phase FeB could lead to cracks during hot rolling and ductility loss (reduction in elongation at break). This can be safely avoided in particular if the ratio % Ti/% B is 0.5-3.75, in particular 1.0-3.75.

The presence of Ti in the flat steel product according to the invention can also improve oxidation resistance and heat resistance. However, if the levels of Ti borides are too high, this would lead to strong solidification when a flat steel product according to the invention is cold formed. Therefore, the upper limit of the Ti levels are 2% by weight, in particular at most 1.5% by weight or 1.1% by weight, and the upper limit of the B levels are 0.60% by weight, in particular at most 0.4% by weight.

Chromium may optionally be present in the steel according to the invention at levels of up to 7% by weight, in particular at least 0.3% by weight, or at least 0.5% by weight, or at least 1.0% by weight in order to lower the brittle-ductile transition temperature and improve overall ductility. The presence of Cr also increases the steel's resistance to low and high temperature corrosion and improves oxidation resistance. There is no increase in these effects at levels higher than 7% by weight, while Cr levels of up to 5% by weight have been found to be particularly effective as far as weighing the cost/benefit, while in practice contents of up to 3% by weight have also proven sufficient to cause improvements of the steel according to the invention related to the addition of Cr.

Carbon, in combination with high Al levels, tends to form embrittling phases (kappa carbides), which reduces hot and cold formability. This would apply in particular if the C levels of a flat steel product according to the invention were higher than 0.15% by weight. Therefore, according to the invention, the lowest possible C levels are sought. However, C entered the steel as an unavoidable production-related impurity, which means that, in practice, levels of at least 0.005% by weight, in particular at least 0.01% by weight, must be expected. In practical tests, it was also found that C levels of up to 0.05% by weight, in particular up to 0.03% by weight, only lead to comparably small impairments of the steel, i.e. they are still acceptable.

The optional addition of manganese in levels of up to 1% by weight can also lower the brittle-ductile transition temperature. In the course of steel production Mn also enters the steel as an unavoidable production-related impurity when Mn is used for deoxidation. Mn helps increase strength, but may deteriorate the corrosion resistance. This is prevented by limiting the maximum Mn levels according to the invention to 2% by weight, in particular max. 1% by weight or max. 0.3% by weight.

Silicon can enter the steel of a flat steel product according to the invention as a deoxidiser in the steel production process, but can also be selectively added to the steel in levels of up to 5% by weight, in particular up to 2% by weight, in order to optimise strength and corrosion resistance, although excessive Si levels can lead to brittle material behaviour. The Si levels of a flat steel product according to the invention are typically at least 0.05% by weight, in particular at least 0.1% by weight.

Phosphorus and sulphur count towards the undesirable but inevitable production-related impurities of a steel according to the invention. The P and S levels should therefore be kept low enough to avoid harmful effects. For this purpose, P levels are limited to max. 0.1% by weight and S levels to max. 0.03% by weight, wherein S levels of max. 0.01% by weight or P levels of max. 0.05% by weight have been found to be particularly advantageous.

Although the optionally present elements niobium, tantalum, tungsten, zirconium and vanadium form strength-enhancing carbides with C in the steel according to the invention, and can contribute to improving heat resistance, at excessively high levels they degrade cold formability and weldability. The latter applies in particular to Nb, Ta and W, which are therefore permitted in the steel according to the invention in levels of at most 0.2% by weight, in particular at most 0.1% by weight. The Zr and V levels are limited toup to 1% by weight in the steel according to the invention, while Zr levels ofup to 0.1% by weight and V levels ofup to 0.5% by weight have been found to be particularly favourable. If the levels are excessive, Zr deteriorates corrosion behaviour, whereas excessive levels of V impair oxidation behaviour. The positive effects of Zr and V can be used especially when at least 0.02% by weight of either Zr or V are present in the steel.

Molybdenum may optionally be added to the steel of a flat steel product according to the invention to improve tensile strength as well as creep resistance and high temperature fatigue strength. Mo can additionally contribute to a fine microstructure by forming fine carbides and complex borides. These positive effects are achieved when the Mo levels are at least 0.2% by weight. However, excessive levels of Mo lead to a deterioration of hot and cold formability. Therefore, the Mo levels of a flat steel product according to the invention are limited to a maximum of 1% by weight, in particular max. 0.7% by weight.

Nickel may optionally be present in the flat steel product according to the invention in levels of up to 2% by weight to improve its strength and ductility, as well as to improve its corrosion resistance. At Ni levels of more than 2% by weight, these effects cease to significantly increase. The positive effects of Ni can be used especially when at least 0.2% by weight, in particular at least 1% by weight, of Ni is present in the steel.

Copper may also optionally be present in the steel according to the invention to improve corrosion resistance. Up to 3% by weight Cu, in particular up to 1% by weight Cu, can be added to the steel for this purpose. The hot formability, weldability and recyclability of a flat steel product according to the invention deteriorate at higher Cu levels. The positive effects of Cu can be used especially when at least 0.2% by weight Cu is present in the steel.

Calcium can be added to the steel during steel production to bind S and prevent clogging during casting of the steel. Optimum effects are achieved in steel compositions according to the invention when the Ca levels are up to 0.015% by weight, in particular max. 0.01% by weight, wherein Ca is safe to use when there is at least 0.001% by weight Ca in the steel.

Rare earth metals “REM” can be added to the steel according to the invention in levels of up to 0.2% by weight, in particular up to 0.05% by weight, to improve oxidation resistance. This effect is achieved in particular if at least 0.001% by weight REM are present in the steel.

Nitrogen is present in the steel according to the invention as an undesirable but usually inevitable production-related impurity. However, in order to avoid harmful factors, the N levels should be kept as low as possible. Limiting the levels of N to max. 0.1% by weight, in particular max. 0.03% by weight, minimises the formation of detrimental Al nitrides, which could degrade mechanical properties and ductility.

Cobalt may optionally be present in the steel according to the invention at levels of up to 1% by weight to increase its heat resistance. This effect is achieved in particular if at least 0.2% by weight Co is present in the steel.

The proportion of TiB₂ in the microstructure of a flat steel product according to the invention is 0.3 to 5% by volume. The presence of such amounts of TiB₂ causes a ductilisation of the Fe₃Al matrix as a result of a significantly increased dislocation density in the vicinity of the TiB₂ particles, and promotes the recrystallisation of the microstructure. At the same time, grain boundary pinning prevents grain coarsening. In order to achieve these effects, at least 0.3% by volume TiB₂ is required in the microstructure, while the effects are particularly reliable when the levels of TiB₂ in the microstructure of the steel according to the invention are at least 0.5% by volume, in particular at least 0.8% by volume. Harmful effects of excessive levels of Ti boride can be reliably prevented by limiting the level of TiB₂ in the microstructure of the flat steel product according to the invention to max. 3% by volume.

Limiting the grain size of the Fe₃Al of the structural matrix to max. 500 μm, in particular max. 100 μm, achieves good strength and ductility at room temperature as well as a good strength at high temperature. Optimally, the average grain size of the Fe₃Al of the structural matrix should be 20-100 μm to ensure sufficient ductility and good creep resistance of the steel at room temperature, wherein average grain sizes of 50 μm have been found in practice to be particularly advantageous.

The effect of the TiB₂ precipitates in the structural matrix of the flat steel product according to the invention can be further optimised if at least 70% of the TiB₂ precipitates in the structural matrix have an average particle diameter of 0.5-10 μm, in particular 0.7-3 μm.

The structural matrix of a flat steel product according to the invention consists of at least 80% by volume of the intermetallic phase Fe₃Al, the aim here being for the matrix to consist as completely as possible, optimally up to 100% by volume, of Fe₃Al. In addition to Fe₃Al, the structural matrix may also contain optional levels of the mixed crystal Fe(Al) or the intermetallic phase FeAl. High levels of at least 80% by volume Fe₃Al are required to set high corrosion resistance, heat resistance, hardness and wear resistance.

To produce a flat steel product according to the invention, a molten steel composed according to the invention in the manner explained above is melted in step a) of the process according to the invention and cast into a precursor in the form of a slab, thin slab or a cast strip in step b). In general, the electric furnace route better suits the operational melting production of a high-alloy steel according to the invention than the classic blast furnace converter route of an integrated smelting plants due to its suitability for liquefaction of high amounts of alloy. The melt can be cast in conventional continuous casting if a suitable casting powder is used. If this proves to be problematic at very high Al levels, it is possible to swap to a near-netshape casting process, such as processes in which the melt is processed into thin slabs that are then processed, without interruption after casting, into hot strips (casting-rolling-process), or into a cast strip, which is then also immediately subjected to a hot rolling process.

For hot rolling (step c)), the respective precursor is brought to the preheating temperature of 1200-1300° C. This can be carried out in a separate heating process or by holding at the relevant temperature from the casting heat. If a separate heating process occurs, it should last for a period of 15-1500 minutes to ensure homogeneous heating. If the temperature or hold time is too low, this is not achieved with the required certainty due to the low thermal conductivity of the steel, which can lead to cracks in the hot strip. A suitable hot rolling start temperature ensures formability, especially in the final passes, and thus avoids high loads on the rolls. The danger of rolling damage as a result of excessive rolling forces can be prevented by choosing a hot rolling start temperature predetermined according to the invention in the range of 1000-1200° C., in particular 1100-1170° C. However, an excessive hot rolling start temperature would result in the material being too weak for hot rolling. This can lead to unwanted deformations during processing and the rolled material sticking to the rollers. According to the invention, the hot rolling end temperature must be at least 850° C. in order to avoid excessive roll forces and to achieve high degrees of formability. At even lower hot rolling temperatures, the required hot strip flatness cannot be guaranteed with the necessary certainty from an operational point of view.

After hot rolling, the hot strip is coiled in step d) at a coiling temperature between room temperature and 750° C. Water or aqueous solutions that ensure a homogeneous cooling over the strip cross-section are particularly suitable as cooling media.

Coiling temperatures of at least 400° C., in particular at least 450° C., have proven particularly useful in practical application, wherein the upper limit of the coiling temperature range can be limited to max. 700° C., in particular max 550° C., to prevent excessive scale formation on the hot strip.

The hot strip obtained after hot rolling has an elongation at break of 2-4% in the tensile test. In order to improve this property, the hot strip can be optionally annealed after coiling at an annealing temperature of 200-1000° C. over an annealing period of 1 to 200 h. This serves to increase ductility at room temperature. A bell annealing process with a peak temperature above 650° C. is suitable for annealing the hot strip. Lower annealing temperatures or holding times show no effect, whereas higher annealing temperatures or holding times can lead to ductility loss due to grain coarsening as a result of coarsening of the Ti boride particles and the Fe₃Al matrix.

The hot strip according to the invention can also be optionally subjected to a pickling treatment with common media, wherein the pickling time must be chosen so that even the stable Al-oxides that appear on the hot strip are eliminated.

In a flat steel product according to the invention, TiB₂ particles are increasingly incorporated in the Fe₃Al intermetallic matrix as a consequence of the high Ti and B levels of the steel from which the flat steel product is made. A flat steel product alloyed according to the invention therefore has high yield points and tensile strengths. At the same time, its density is greatly reduced compared to conventional steels of the same strength class. The typical density of steels according to the invention is in the range of 6.2-6.7 g/cm³ and is on average typically 6.4 g/cm³. This results in a high strength/density ratio compared to other heat-resistant materials.

By selecting the rolling parameters according to the invention, the BDTT value (brittle-ductile transition) can be lowered to surprisingly low temperatures of about 75-100° C.

Above this temperature, the elongation at break increases with increasing temperature and reaches extremely high values at 650° C. Due to the ductility that increases with an increase in temperature, it is feasible to produce components using preheated sheets or conventional hot forming.

Typical hot yield points of flat steel products according to the invention are at 650° C. with approx. 130-170 MPa in the range of conventional ferritic Cr steels, such as the steel standardised under material number 1.4512 (hot yield point approx. 70 MPa) and the steel designed for high heat resistance, standardised under material number 1.4509 (hot yield point approx. 150 MPa). At temperatures of at least 700° C., the tensile strength of the flat steel product according to the invention is still regularly at at least 100 MPa.

Due to their property profiles, flat steel products produced according to the invention are particularly suitable for the production of, in particular, heat-resistant components for plant construction (e.g. heavy plate), gas turbines, offshore installations and, in particular, heat-resistant components for the automotive industry, in particular exhaust systems or turbocharger housings (hot strip). Other preferred uses are conceivable in the low temperature range (e.g. biogas plants, brake discs, vehicle underbodies).

The invention is explained in greater detail below using example embodiments.

Sixty kg of each of the alloys A-F indicated in Table 1 were melted with argon in a vacuum induction furnace and poured into moulds measuring 250×150×500 mm. After solidification, the ingots obtained were preheated to 1200° C., then rolled down to 45 mm on a two-high reversing stand and divided into six blooms with a height of 40 mm each. The resulting blooms were heated to a preheat temperature of 1200° C. over a preheat time of 180 minutes each.

The heated blooms were each hot rolled into a hot strip with a thickness of 3 mm in a conventional manner at a hot rolling end temperature HRET from a hot rolling start temperature HRST.

The obtained hot strips were cooled from the respective hot rolling end temperature HRET to the respective coiling temperature CT and wound into a coil at that temperature.

The HRST, HRET and CT parameters for the different samples A1-F3 are shown in Table 2.

The following was determined for samples A1-F3: the mechanical properties yield strength Rp0.2, tensile strength Rm and elongation A50 at room temperature (see Table 3), and also for some selected samples, the mechanical properties yield strength Rp0.2, tensile strength Rm and elongation at break A at 650° C. (see Table 4), as well as the texture properties “Grain size of the matrix”, “Matrix” and “TIB₂ content in the microstructure” (see Table 5) and the brittle transition temperature BDTT (see Table 6).

The mechanical properties were determined in the tensile test according to DIN EN 10002, whereas the brittle-ductile transition temperature was determined in the four-point bending test. The four-point bending tests were performed on 3×6×18 mm³ sized samples between room temperature and 500° C. The samples were wet sanded longitudinally with 1000 grit sandpaper prior to the start of the test. The experiments were carried out with a deformation rate of phi=1×10⁻⁴ s⁻¹ in air. This is the standard method for determining the brittle-ductile transition temperature for intermetallic phases (see D. Risanti et al. “Dependence oft he brittle-to-ductile transition temperature BDTT on the Al content of the Fe—Al alloys”; Intermetallics, 13(12), (2005) 1337-1342). The grain size of the matrix was determined in the line-cut method according to DIN ISO 643. TiB₂ particle size and volume fraction were determined according to ASTM E 1245.

It was found that the alloys A-F could easily be rolled using industrial conditions on a laboratory scale.

The tests have thus confirmed that the tensile strengths Rm of flat steel products according to the invention at room temperature typically have 550-700 MPa and yield points Rp0.2 of 400-650 MPa at an elongation A50 of typically 2-5%. The tensile strength could be increased in particular if the pre-roll and finishing roll occurred in different rolling directions.

The Vickers hardness HV5 typically varies between 335 and 370 in the flat steel products according to the invention.

The hot yield point 60.2 (measured transversely to the rolling direction according to DIN EN 10002) at 650° C. is typically 120±170 MPa.

In the 4-point bending test it was found that the sheets do not have a pronounced brittle-ductile transition temperature of 75-100° C. They are fully ductile even at 100° C. This is an improvement of at least 150° C. compared to the cast material, and is due to the microstructure being refined by rolling. Ductility can be increased by hot strip annealing of the type described above.

TABLE 1 All data in % by weight, residual iron and other production-related unavoidable impurities Nb, W, Ta, Steel Al B Ti Cr Mn Si Ni Cu C N P S Zr, V each Mo Mo REM Co Ti/B A 14.7 0.2 0.62 0.1 0.45 0.35 <0.1 <0.1 0.018 0.009 0.013 0.001 <0.01 <0.01 0.0015 <0.001 <0.001 3.10 B 14.8 0.31 0.98 0.1 0.38 0.6 <0.1 <0.1 0.017 0.007 0.015 0.001 <0.01 <0.01 0.0014 <0.001 <0.001 3.16 C 18.7 0.2 0.65 0.1 0.38 0.54 <0.1 <0.1 0.02 0.010 0.014 0.001 <0.01 <0.01 0.0013 <0.001 <0.001 3.25 D 14.2 0.2 0.67 0.51 0.41 0.42 <0.1 <0.1 0.02 0.007 0.014 0.001 <0.01 <0.01 0.0014 <0.001 <0.001 3.35 E 13.9 0.29 0.89 1.86 0.51 0.53 <0.1 <0.1 0.018 0.008 0.013 0.001 <0.01 <0.01 0.0015 <0.001 <0.001 3.07 F 19.2 0.21 0.75 0.67 0.47 0.58 <0.1 <0.1 0.019 0.009 0.014 0.001 <0.01 <0.01 0.0014 <0.001 <0.001 3.57

TABLE 2 HRST HRET CT Sample Steel [° C.] [° C.] [° C.] A1 A 1140 870 500 A2 1150 970 700 A3 1150 940 RT B1 B 1130 940 500 B2 1160 970 700 B3 1170 950 RT C1 C 1150 940 500 C2 1140 950 700 C3 1130 930 RT D1 D 1160 880 500 D2 1150 920 RT E1 E 1170 860 500 E2 1160 950 700 E3 1140 920 RT F1 F 1150 930 500 F2 1170 940 700 F3 1130 920 RT

TABLE 3 Rp0.2 Rm A50 Sample [MPa] [MPa] [%] A1 418 589 2.7 A2 419 574 2.9 A3 476 601 2.4 B1 442 631 2.2 B2 425 614 2.5 B3 598 675 2 C1 432 601 2.1 C2 417 589 2.4 C3 629 689 2 D1 425 596 4.1 D2 445 554 3.9 E1 455 641 3.7 E2 435 619 3.9 E3 561 629 3.4 F1 439 615 2.7 F2 419 595 2.8 F3 594 654 2.4

TABLE 4 Rp0.2 Rm A Sample [MPa] [MPa] [%] A3 128 128 69 B3 145 146 85 C3 168 170 57 D2 125 126 68 E3 139 140 88 F3 161 163 62

TABLE 5 Average grain size of the matrix [μm] on TiB₂ content Sample longitudinal sample Matrix [% by volume] A3 46 Fe₃Al (type DO₃) 0.9 B3 53 Fe₃Al (type DO₃) 1.4 C3 55 Fe₃Al (type DO₃) 1 D2 48 Fe₃Al (type DO₃) 1 E3 64 Fe₃Al (type DO₃) 1.4 F3 46 Fe₃Al (type DO₃) 1

TABLE 6 Sample State BDTT Notes A3, B3, C3 Hot strip 100° C. According to the invention D2, E3, F3 Hot strip  75° C. According to the invention A + C Cast 250° C. Comparison B, C, E, F Cast 350° C. Comparison 

1. A flat steel product made from a steel that comprises (in % by weight): Al: 12-20% Ti: 0.2-2% B: 0.1-0.6%, as well as optionally one or more elements selected from the group consisting of Cr, C, Mn, Si, Nb, Ta, W, Zr, V, Mo, Ni, Cu, Ca, rare earth metals, and Co at the following levels: N: up to 0.1% Cr: up to 7% C: up to 0.15% Mn: up to 2% Si: 0.05-5% Nb, Ta, W: up to 0.2% in total Zr: up to 1% V: up to 1% Mo: up to 1% Ni: up to 2% Cu: up to 3% Ca: up to 0.015% Rare earth metals: up to 0.2% Co: up to 1% remainder iron and unavoidable impurities, wherein S levels of up to 0.03% by weight and P levels of up to 0.1% by weight can be attributable to the unavoidable impurities, and wherein the ratio % Ti/% B formed from the Ti levels % Ti and the B levels % B of the steel is 0.33≤% Ti/% B≤3.75 and a microstructure of the flat steel product comprises 0.3-5% by volume TiB₂ precipitates that are embedded in a structural matrix made up of at least 80% by volume Fe₃Al.
 2. The flat steel product according to claim 1, wherein the % Ti/% B ratio is 0.5≤% Ti/% B≤3.75,
 3. The flat steel product according to claim 2, wherein the % Ti/% B ratio is 1.0≤% Ti/% B≤3.75.
 4. The flat steel product according to claim 1, wherein a grain size of the Fe₃Al in the structural matrix is a maximum of 500 μm.
 5. Thee flat steel product according to claim 4, wherein the grain size of the Fe₃Al in the structural matrix is maximum of 100 μm.
 6. The flat steel product according to claim 1, wherein at least 70% of the TiB₂ precipitates in the structural matrix have a mean particle diameter of 0.5-10 μm.
 7. The flat steel product according to claim 1, wherein the sum of Nb, Ta, and W is up to 0.1% by weight.
 8. The flat steel product according to claim 1, wherein Cr is at least 0.3% by weight.
 9. The flat steel product according to claim 1, wherein the microstructure of the flat steel product comprises at least 0.5% by volume TiB₂ precipitates.
 10. The flat steel product according to claim 1, wherein the microstructure of the flat steel product comprises a maximum of 3% by volume TiB₂ precipitates.
 11. A method for producing the flat steel product according to claim 1, comprising: a) melting a steel that comprises (in % by weight): Al: 12-20% Ti: 0.2-2% B: 0.10-0.6%, as well as optionally one or more elements selected from the group consisting of Cr, C, Mn, Si, Nb, Ta, W, Zr, V, Mo, Ni, Cu, Ca, rare earth metals, and Co at the following levels: N: up to 0.1% Cr: up to 7% C: up to 0.15% Mn: up to 2% Si: 0.05-5% Nb, Ta, W: up to 0.2% in total Zr: up to 1% V: up to 1% Mo: up to 1% Ni: up to 2% Cu: up to 3% Ca: up to 0.015% Rare earth metals: up to 0.2% Co: up to 1% remainder iron and unavoidable impurities, wherein S levels of up to 0.03% by weight and P levels of up to 0.1% by weight can be attributed to the unavoidable impurities, and wherein the ratio % Ti/% B formed from the Ti levels % Ti and the B levels % B of the steel is 0.33≤% Ti/% B≤3.75; b) casting the steel melt into a precursor in a form of a slab, thin slab or cast strip; c) hot rolling the precursor into a hot-rolled hot strip, wherein the precursor has a hot rolling start temperature of 1000-1300° C. at the start of the hot rolling process, and has a hot rolling end temperature of at least 850° C.; and d) coiling the hot strip at a coiling temperature between room temperature and 750° C.
 12. The method according to claim 11, wherein the hot strip obtained after coiling (step d)) is annealed at an annealing temperature of 200-1000° C. over an annealing period of 1-200 hours.
 13. The method according to claim 11, wherein the precursor is heated to the hot rolling start temperature over a heating time of 15-1500 min between steps b) and c).
 14. The method according to claim 11, wherein the coiling temperature is at least 400° C.
 15. (canceled)
 16. Components for plant construction comprising the flat steel product of claim
 1. 17. Components for gas turbines comprising the flat steel product of claim
 1. 18. Heat-resistant components for the automotive industry comprising the flat steel product of claim
 1. 19. Components for plants that are operated in the low-temperature range comprising the flat steel product of claim
 1. 20. Hot formed component comprising the flat steel product of claim
 1. 